Methods for making a porous nuclear fuel element

ABSTRACT

Porous nuclear fuel elements for use in advanced high temperature gas-cooled nuclear reactors (HTGR&#39;s), and to processes for fabricating them. Advanced uranium bi-carbide, uranium tri-carbide and uranium carbonitride nuclear fuels can be used. These fuels have high melting temperatures, high thermal conductivity, and high resistance to erosion by hot hydrogen gas. Tri-carbide fuels, such as (U,Zr,Nb)C, can be fabricated using chemical vapor infiltration (CVI) to simultaneously deposit each of the three separate carbides, e.g., UC, ZrC, and NbC in a single CVI step. By using CVI, the nuclear fuel may be deposited inside of a highly porous skeletal structure made of, for example, reticulated vitreous carbon foam.

CROSS-REFERENCE TO RELATED APPLICATIONS

This application is a divisional application of U.S. Ser. No.12/850,752, Porous Nuclear Fuel Element for High-Temperature Gas CooledNuclear Reactors, by Youchison et al., filed Aug. 5, 2010, issued onSep. 3, 2013 as U.S. Pat. No. 8,526,566, which is a divisionalapplication of U.S. Ser. No. 11/435,412, Methods for ManufacturingPorous Nuclear Fuel Elements for High-Temperature Gas-Cooled NuclearReactors, by Youchison et al., filed on May 17, 2006, issued on Feb. 23,2010 as U.S. Pat. No. 7,666,463, which is a divisional application ofU.S. Ser. No. 10/881,873, Porous Nuclear Fuel Element forHigh-Temperature Gas Cooled Nuclear Reactors, by Youchison et al., filedJun. 29, 2004, issued on Mar. 1, 2011 as U.S. Pat. No. 7,899,146, all ofwhich are incorporated herein by reference.

FEDERALLY SPONSORED RESEARCH

The United States Government has rights in this invention pursuant toDepartment of Energy Contract No. DE-AC04-94AL85000 with SandiaCorporation.

BACKGROUND OF THE INVENTION

The present invention relates generally to porous nuclear fuel elementsfor use in advanced high temperature gas-cooled nuclear reactors(HTGR's), and to methods for fabricating same.

Exploring the solar system while maintaining reasonable interplanetarytravel times will require increases in spacecraft velocities of one totwo orders of magnitude over currently achievable levels. The “packagedenergy” for space propulsion systems required to achieve this goal,expressed as a combination of specific impulse (I_(sp)) and massfraction, must increase by an order of magnitude over current chemicalrocket propulsion. The primary problem is that even advanced chemicalrocket systems limit maneuverability and destinations. Chemical systemswill not be able to extend human space exploration much beyond the moonor Mars. One solution to this problem is to utilize more energeticfuels. Therefore, NASA's goal is to develop fission propulsion to enablerapid and affordable access to any point in the solar system.

Fission power enables a new propulsion growth path with two options:nuclear thermal propulsion (NTP), in which gas is heated and expandedthrough a nozzle; or nuclear electric propulsion (NEP), in which nuclearpower is converted to electric power for an advanced propulsion system,such as a plasma thruster. The I_(sp) for NTP systems can increase by afactor of 10 to 100, and the specific energy can improve by a factor of1,000,000 over conventional rockets. NTP systems have high thrust withmoderate I_(sp) (>800 sec), whereas NEP systems have lower thrust withextended I_(sp) (>5000 sec) while providing auxiliary payload power.When compared to chemical power, fission power provides more energy forinstruments, communications and higher data processing rates; more timefor extended exploration; and more adaptability to changing missionrequirements.

Recent conceptual designs, e.g., SAFE-100 and SAFE-300, are based onusing uranium oxide (UO₂) nuclear fuel elements. A 30-kW resistivelyheated prototype module, SAFE-30, underwent testing to verify heattransfer through an annular core geometry utilizing heat pipes and a350-W Stirling engine power converter. Such designs clearly havepotential for lower thrust, long-I_(sp) NEP systems, but are limited bytheir maximum operating temperature when high-thrust, shorter-pulse NTPsystems or bimodal NTP/NEP systems would best meet mission requirements.

The efficiency of the rocket increases if the temperature difference ΔTbetween the fuel and propellant is minimized. This dramatically improvesthrust-to-mass ratio, reduces the amount of propellant required(tankage) for NTP systems, and can improve total available electricpower for NEP systems when using a high-efficiency Brayton thermal cycleor advanced Stirling engine. Current designs for high-temperaturegas-cooled reactor fuel elements, such as annual rods or pebble beds,cannot operate at extremely high temperatures and, thus, have limitedefficiency.

U.S. Pat. No. 5,094,804 to Schweitzer and U.S. Pat. No. 4,659,911 toBingham et al. are representative of older types of nuclear fuelelements.

Helium gas-cooled reactors have been operated in the USA usingpebble-bed core designs with spherical fuel elements (60 mm diameter“pebbles”) made of TRISO fuel particles embedded in a graphite matrix.Each TRISO fuel particle is a microsphere (0.9 mm diameter) comprising akernal/core of fissile material (UO₂ or UC₂) coated by multiple layersof protective barrier materials, e.g., a porous carbon buffer layer,followed by pyrocarbon, followed by silicon carbide, followed by anexternal coating of pyrocarbon. A typical reactor core might contain11,000 fuel pebbles.

A new project, Prometheus, will enable advanced human exploration of thesolar system, including Mars, and beyond. It specifically calls for thedevelopment of new nuclear fuels and components that are capable ofextremely high temperature and very efficient operation (>925 sec I_(sp)for >1 hr). An attractive approach would be to use high-efficiency,gas-permeable, porous nuclear fuel elements for high-temperature,gas-cooled (e.g., hydrogen, helium) space reactors. The porous fuelelements can operate at extremely high temperatures when usingrefractory materials with low neutron absorption cross-sections, and canhave much higher heat transfer coefficients due to their very largesurface-to-volume ratio and extended surface areas for transferringheat, which allows for greater power densities. The interior surfaces ofthe porous body could be coated with a thin layer of the nuclear fuel,so that the (hot) hydrogen or helium gas would easily flow through theinterconnected open porosity to efficiently exchange heat generated bythe nuclear fuel to the gas coolant. Such a porous nuclear fuel elementincorporating an enriched uranium (or plutonium, thorium, americium) bi-or tri-carbide fuel could have extremely high surface area andstiffness, low density (light-weight), extremely high melting point andexcellent thermal conductivity; would not degrade in hydrogen at 3000 K,would not clog; and would retain its structural integrity at hightemperatures. The dispersed fissile material could be an integral partof a high-conductivity matrix. Thus, a greatly reduced temperaturedifference between the center of the fuel and the gas coolant (i.e.,propellant) temperature could exist, which would allow the gastemperature to be much higher, yielding the high specific impulserequired to sustain interplanetary exploration.

Highly porous (e.g., 90% porous) metal carbide foam structures (withoutnuclear fuel) have been successfully fabricated in the aerospaceindustry by Ultramet, Inc. through chemical vapor deposition of one ormore layers of a refractory metal carbide, for example, ZrC or NbC, on aporous foam skeleton made of, for example, reticulated vitreous carbon(RVC). These metal carbide foams have been used as thermal protectionsystems, actively cooled structures/heat exchangers, flash and blastsuppressors, and lightweight mirror substrates. The interconnected opencell geometry and tortuous flow path provides excellent heat exchangeproperties, excellent particulate filtration, with a correspondingly lowmass. Just about any material that can be deposited by CVD/CVItechniques can be used to make a porous structure by depositing themonto a skeletal structure (e.g., RVC foam). Examples of suitablematerials that can be deposited by CVD/CVI include, but are not limitedto: Zr, Nb, Mo, Hf, Ta, W, Re, TiC, TaC, ZrC, SiC, HfC, BeC₂, B₄C, NbC,GdC, HfB₂, ZrB₂, Si₃N₄, TiO₂, BeO, SiO₂, ZrO₂, HfO₂, Y₂O₃, Al₂O₃, Sc₂O₃,and Ta₂O₅. Foam structures made of NbC and/or ZrC deposited on a RVCmatrix) have several important advantages over bulk high temperaturematerials, such as low overall density, lack of degradation in hothydrogen at 2700 C (where they also retain structural integrity), andminimal neutron cross-section (i.e., reduced parasitic neutronabsorbtion).

Solid solution, mixed carbide fuels, such as uranium carbide (UC, UC₂)and uranium bi-carbide fuels (U,Zr)C were studied in the 1970's fornuclear thermal propulsion of spacecraft as part of the Rover/NERVAprogram in the USA, and in similar programs in the former Soviet Union.Fuel elements designs included dispersions of small particles of UO₂ orUC₂ in solid graphite blocks; and a composite design made ofsolid-solution (U,Zr)C dispersed inside of graphite. Both fuel typeswere protected by a NbC or ZrC fission product barrier coating. Theseprotective coatings were needed to protect against unacceptable massloss due to the high chemical reactivity of free carbon with the flowinghot hydrogen propellant; and due to mis-matches in thermal expansioncoefficient between the graphite matrix and the NbC or ZrC coatings. Atthe time the program was cancelled in 1973, there had been an evolutionin thinking towards considering an all-carbide, solid-solution uraniumbi-carbide fuel, e.g., (U,Zr)C or (U,Nb)C, because of their expectedhigh resistance to erosion from exposure to hot hydrogen gas. However,the uranium bi-carbide fuels were never infiltrated into a porous matrixstructure to make a porous fuel element. In addition, single-phase,solid-solution uranium tri-carbide fuels, such as (U,Zr,Nb)C, were neverseriously considered because they had not been synthesized or fabricatedat that point in time. It has only been in recent years that uraniumtri-carbide fuels have been successfully fabricated, and their basicproperties measured.

Advanced uranium (or plutonium) tri-carbide fuels have been proposed fornuclear thermal propulsion (NTP) applications because of their expectedlonger life and higher operating temperature; due to their high meltingtemperature, high thermal conductivity, and improved resistance to hothydrogen corrosion. Recently, high density (e.g., 95%), solid solutionmixed uranium/refractory metal tri-carbide fuels have been manufacturedusing a high-temperature liquid-phase sintering technique developed atthe University of Florida (see T. W. Knight & A. Anghaie, “Processingand fabrication of mixed uranium/refractory metal carbide fuels withliquid-phase sintering”, Journal of Nuclear Materials 306 (2002) p.54-60.)

These uranium tri-carbide fuels, such as (U, Zr, X) C with X=Nb, Ta, Hf,or W, exhibit high melting temperatures (greater than 3400 C) foruranium metal mole fractions of 10% or less. This melting point isalmost 1000 C higher than pure UC; and UC₂ has an even lower meltingpoint. Also, UC₂ erodes much more quickly in hot hydrogen than UC. Mixedphases of UC and UC₂ lead to eutectic melting at a temperature 500 Clower than UC. For the tri-carbide fuels, uranium fractions greater than10% lower the melting temperature and lead to greater uranium mass lossfrom either interactions with the flowing hot hydrogen propellant orvaporization from the fuel element surface, especially near the bottomof the reactor core where fuel surface and hydrogen exit temperaturesmight be expected to exceed 2500 C. The high solid-phase solubility ofUC with the refractory metal carbides (ZrC, NbC, TaC, etc.) permits alarge degree of flexibility in designing uranium bi-carbide ortri-carbide nuclear fuel elements. The carbon-to-metal (C/M) ratio canbe less than 0.95 in order to maintain high melting point. When theuranium concentration is less than about 10% mole fraction, then themelting point of the tri-carbide can be as high as 3400 C for a solidsolution.

Hydrogen testing on these solid solution uranium tri-carbide alloy fuelshas been performed, with little erosion observed at 2700-2800 C.Additionally, the thermal conductivities of these solid-solution uraniumtri-carbide fuels is much higher than conventional uranium oxide nuclearfuels (by factors of 10-20 times higher). These superior physicalproperties improves overall reactor efficiency and reduces system costby allowing higher operating temperatures, reducing the amount ofnuclear fuel (owing to small, more compact cores), and reducingpropellant requirements (including reduced refrigeration costs). Hence,solid solution uranium tri-carbide fuels are preferred over uraniumsingle-carbide and bi-carbide fuels.

Unfortunately, the method of fabrication described above for producingthese advanced uranium tri-carbide fuels (i.e., liquid-phase sintering),cannot be used to deposit thin coatings of nuclear fuel onto exposedinterior surfaces of a highly porous foam skeleton made of, e.g.,reticulated vitreous or glassy carbon. The sintering step would quicklydestroy the thin, interconnected structural ligaments. Therefore, theprevious process of liquid-phase sintering cannot be used to make aporous, gas-permeable fuel element made of uranium tri-carbide fuel.

What is needed, then, is a high-efficiency, gas-permeable, porousnuclear fuel element for use in high temperature gas-cooled nuclearreactors (HTGR's), and a process for fabricating them, which utilizesadvanced uranium bi-carbide, uranium tri-carbide, and uraniumcarbonitride nuclear fuels having higher melting temperatures, higherthermal conductivity, and improved resistance to corrosion from hothydrogen gas.

Against this background, the present invention was developed.

SUMMARY OF THE INVENTION

The present invention relates to porous nuclear fuel elements for use inadvanced high temperature gas-cooled nuclear reactors (HTGR's), and toprocesses for fabricating them. Advanced uranium bi-carbide, uraniumtri-carbide and uranium carbonitride nuclear fuels can be used. Thesefuels have high melting temperatures, high thermal conductivity, andhigh resistance to erosion by hot hydrogen gas. Tri-carbide fuels, suchas (U,Zr,Nb)C, can be fabricated using chemical vapor infiltration (CVI)to simultaneously deposit each of the three separate carbides, e.g., UC,ZrC, and NbC in a single CVI step. The nuclear fuel may be depositedusing CVI inside of a highly porous skeletal structure made of, forexample, a reticulated vitreous carbon foam.

BRIEF DESCRIPTION OF THE DRAWINGS

The accompanying drawings, which are incorporated in and form part ofthe specification, illustrate various examples of the present inventionand, together with the detailed description, serve to explain theprinciples of the invention.

FIG. 1 illustrates a schematic side view of an example of a porousnuclear fuel element, according to the present invention.

FIG. 2A is a SEM micrograph magnified 15× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton before being infiltrated withnuclear fuel. The average pore diameter is about 800 microns and thetotal porosity is about 97%.

FIG. 2B is a SEM micrograph magnified 45× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton before being infiltrated withnuclear fuel. The average pore diameter is about 800 microns and thetotal porosity is about 97%.

FIG. 2C are SEM micrographs of a reticulated vitreous carbon (RVC) foamskeleton, illustrating the meanings of “cell structure”, “windowdiameter”, “pore diameter”, and “pores-per-inch (PPI)”.

FIG. 3 illustrates a schematic view of an example of a porous body withinternal structure coated by nuclear fuel, according to the presentinvention.

FIG. 4 illustrates a schematic cross-section view, Section A-A, of anexample of a ligament of a porous body coated with nuclear fuel and anouter barrier coating, according to the present invention.

FIG. 5 illustrates a schematic cross-section view, Section A-A, of anexample of a ligament of a porous body, with the ligament coated by alayer of nuclear fuel and two layers refractory metal carbides,according to the present invention.

FIG. 6 is a SEM micrograph magnified 15× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton before being infiltrated withnuclear fuel. The RVC skeleton has been compressed in one direction tocreate elongated pores aligned along an axis.

FIG. 7A is a SEM micrograph magnified 15× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after infiltration with TaC—NbC—ZrC.

FIG. 7B is a SEM micrograph magnified 45× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after infiltration with TaC—NbC—ZrC.

FIG. 8A is a SEM micrograph magnified 250× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after infiltration with individuallayers of TaC/NbC/ZrC. The outer surface of ZrC appears to be continuousand crack-free.

FIG. 8B is an optical micrograph of a polished cross-section of alayered tri-carbide foam sample, magnified 40×. The individual layers ofthe RVC foam skeleton, TaC, NbC, and ZrC can be seen.

FIG. 8C is a SEM micrograph magnified 400× of a polished cross-sectionof a layered tri-carbide foam sample. The individual layers of TaC, NbC,and ZrC can be seen. However, because this is a backscatteredsecondary-electron image (BSE) the inside core of RVC foam skeleton cannot be seen (since carbon is a low-Z material, it appears black, as doesthe central void). The light-colored band is the NbC layer.

FIG. 9A is a SEM micrograph magnified 27× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after being infiltrated with layersof TaC/NbC/ZrC to a total of 10% vol. density. TaC was used in thissample as a substitute for UC.

FIG. 9B is a SEM micrograph magnified 400× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after being infiltrated with layersof TaC/NbC/ZrC to a total of 10% vol. density. Because this is abackscattered primary-electron image, the inside core of RVC foamskeleton can be easily identified.

FIG. 10A is a SEM micrograph magnified 27× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after being infiltrated with layersof TaC/NbC/ZrC to a total of 30% vol. density.

FIG. 10B is a SEM micrograph magnified 400× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after being infiltrated with layersof TaC/NbC/ZrC to a total of 30% vol. density.

FIG. 11A illustrates a schematic view of an example of a porous nuclearfuel element, where the interconnecting ligaments of the porous body aremade of nuclear fuel, according to the present invention.

FIG. 11B illustrates a schematic cross-section view, Section B-B, of anexample of a hollow, triangular-shaped ligament of a porous body,wherein the ligament is made of nuclear fuel, according to the presentinvention

FIG. 12 illustrates a schematic isometric view of an example of a porousnuclear fuel element comprising porous nuclear fuel encased within acylinder of metal cladding, according to the present invention.

FIG. 13 illustrates a schematic cross-section view of an example of anannular porous nuclear fuel element in the shape of a thick-walledcylinder, having a closed end, where gas cooling enters through the opencentral core and exits by flowing radially inwards or outwards throughthe porous fuel.

DETAILED DESCRIPTION OF THE INVENTION

Herein, when we refer to a tri-carbide nuclear fuel, e.g., (U,Zr,Nb)C,we broadly define this to include hyper- and hypo-stoichimetriccompositions, such as (U_(0.1)Zr_(0.77)N_(0.13)) C_(0.95), in additionto stoichimetric compositions. Additionally, when we refer to uranium,we broadly define this as including any fissile element, includingenriched uranium, americium, plutonium, and mixtures thereof, unlessotherwise specifically stated.

FIG. 1 illustrates a schematic side view of an example of a porousnuclear fuel element, according to the present invention. Fuel element10 comprises a porous body 12 comprising internal structure 14 andinterconnected pores 16, 16′, etc. In this example, internal structure14 is illustrated as a fibrous body comprising a tangled web or wovenmesh of fibers or ligaments. Nuclear fuel 18 is disposed on the externalsurfaces of internal structure 14. Fuel 18 comprises one or more alloysselected from the group consisting of solid solution uraniumbi-carbides, solid solution uranium tri-carbides, and solid solutionuranium carbonitrides. The solid solution uranium bi-carbides maycomprise (U,Zr)C or (U,Nb)C, or combinations thereof. The solid solutionuranium tri-carbides may comprise (U,Zr,Nb)C, (U,Zr,Ta)C, (U,Zr,Hf)C, or(U,Zr,W)C, or combinations thereof. The solid solution uraniumcarbonitrides may comprise (U,Zr)CN or (U,Ta)CN, or combinationsthereof. Nuclear fuel 18 may comprise uranium tri-carbide, (U_(W)Zr_(X)Nb_(Y))C_(Z), where 0.04<W<0.12, 0.45<X<0.9, 0<Y<0.45, and0.92<Z<1.0. Alternatively, nuclear fuel 18 may comprise uraniumtri-carbide nuclear fuel having a stoichiometry of about(U_(0.1)Zr_(0.77) Nb_(0.13))C_(0.95). Alternatively, nuclear fuel 18 maycomprise uranium tri-carbide nuclear fuel having a stoichiometry ofabout (U_(0.1) Zr_(0.675) Nb_(0.225))C. Nuclear fuel 18 may comprise athin coating fabricated by simultaneously vapor depositing all of theelements that make up the layer. For example, a solid solution uraniumtri-carbide layer may be fabricated by simultaneously CVI depositing thethree separate carbides UC, NbC, and ZrC at the same time (i.e., “vaporphase alloying”). The solid solution uranium tri-carbide layer 18 formedin this way is a true single-phase alloy of the three individualcarbides, intimately mixed at the molecular scale.

Nuclear fuel element 10, including internal structure 14 coated bynuclear fuel 18, may have a total porosity greater than about 70%.Alternatively, the total porosity may be greater than about 70% and lessthan about 90%. Alternatively, the total porosity may be greater thanabout 77% and less than about 85%. One example of an optimum totalporosity is about 77%, which provides a good balance between heattransfer and pressure drop. This optimum design provides just enoughfissile material to be critical, without resorting to excessively largereactor sizes or using weapon's grade enrichments. The fuel matrix isadaptable to both thermal and fast reactors by inclusion or deletion ofmoderator material (e.g. ZrH or C). Having a thin thickness of thenuclear fuel allows for a high total porosity to maintain the heattransfer efficiency and to keep the temperature difference between theligament (fuel) centerline and the coolant (e.g., helium or hydrogen)bulk temperature as low as possible.

FIG. 2A is a SEM micrograph magnified 15× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton before being infiltrated withnuclear fuel. The average pore diameter is about 800 microns and thetotal porosity is about 97%.

FIG. 2B is a SEM micrograph magnified 45× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton before being infiltrated withnuclear fuel. The average pore diameter is about 800 microns and thetotal porosity is about 97%.

FIG. 2C are SEM micrographs of a reticulated vitreous carbon (RVC) foamskeleton, illustrating the meanings of “cell structure”, “windowdiameter”, “pore diameter”, and “pores-per-inch (PPI)”. For example, thewindow diameter of the RVC foam skeleton in FIGS. 2A and 2B is about 250microns; the average ligament thickness is about 75 microns, with atriangular cross-section; and a specific surface area of about 3700m²/m³ at 97% porosity. Essentially 100% of the porosity in this sample(and for RVC foams in general) comprises interconnected pores.

A first example of a process for making a metal or ceramic coated RVCfoam is as follows. First, commercially available polyurethane foam ispurchased in the desired pore size. Then, the polyurethane foam isinfiltrated with a carbon-bearing resin and pyrolyzed to form a porous,open-celled material composed essentially of vitreous (glassy) carbon,which is called a reticulated vitreous carbon (RVC) foam skeleton. Someshrinkage occurs at this step, but the original pore structure of thepolyurethane foam is retained during conversion to RVC. The RVC foam hasan extremely high void volume (97%), combined with self-supportingrigidity. Pore densities from 3 to 100 pores per linear inch (ppi) arereadily available, and higher ppi foams can be made by compressing100-ppi material prior to pyrolysis in one, two, or three dimensions(see, for example, FIG. 6). Compression or stretching can also be usedto create directional properties (e.g., strength, pressure drop, etc.).Additionally, the RVC foam skeleton can be machined to near finaldimensions prior to vapor infiltration.

High ppi (i.e., hundreds of ppi) compressed carbon foam may be used asthe skeletal structure as a means of increasing the surface area andheat transfer of the nuclear fuel. 65 ppi foam was selected for initialdevelopment, but foams up to 130 ppi are readily produced through resininfiltration of pyrolysis of polyurethane foam, yield vitreous carbon.By compression of foam prior to the conversion to carbon, foams inexcess of 1000 ppi have been fabricated. Compressed foams may beanisotropic structures with directional fluid flow, thermal, andmechanical properties, which may be tailored.

Then, the RVC foam is infiltrated with the desired nuclear fuel to thedesired overall density by using chemical vapor infiltration (CVI), orsome other vapor, liquid, or physical deposition process. Typicalinfiltration levels, depending on the application, fall in the 10-30 vol% range (added to the 3 vol % dense RVC skeleton). At this stage, thethermal and mechanical properties of the foam are mostly dictated by theinfiltrated material. The original RVC foam skeleton has littleinfluence on the final foam properties, and can often be removed throughreaction with hydrogen or oxygen, depending on the particular materialthat was infiltrated.

Chemical vapor infiltration (CVI), a variation of the chemical vapordeposition (CVD) process, is used primarily for depositing materialinside of the porous foam, felt, mesh, or fibrous preform. The vapordeposition process is an extremely versatile and relatively inexpensivemethod of molecular-forming materials that are difficult to machine orotherwise produce by conventional processes. CVI relies on thedecomposition of a gaseous precursor, flowed over (in the case of CVD),or through (in the case of CVI) a heated substrate, with subsequentcondensation from the vapor state to form a solid deposit on thesubstrate. Benefits of CVD/CVI include the ability to produce depositsof controlled density, thickness, orientation, and composition. Impuritylevels are typically less than 0.1%, with densities up to 99.9% beingachievable. In addition, CVD/CVI coating processes exhibit the greatestthrowing power, or ability to uniformly deposit on intricately shaped ortextured substrates. Vapor deposition of ceramic materials inside ofporous substrates possesses distinct advantages over other methods, suchas slurry impregnation, in that precise control of coating thickness,homogeneity, and density can be achieved using CVI.

Perhaps the greatest benefit of CVD/CVI is that a wide variety ofmaterials can be deposited at temperatures that are 10% to 50% of themelting point of the coating material itself, which eliminates the needto perform liquid-phase infiltration at high temperatures. Inpreparation for infiltration, the RVC foam substrate/skeleton can easilybe machined to near final dimensions, while accounting for minordimensional changes that occur during infiltration.

In the CVI process, reactant gases (typically metal chlorides orfluorides containing the desired coating material(s)) are flowed througha heated substrate (e.g., RVC foam). The compound(s) within the reactantgas stream react near the heated ligament surfaces to form a continuous,uniform coating. For example, NbC is deposited at 1000-1200° C. viareaction of niobium pentachloride (NbCl₅) with methane (CH₄) andhydrogen (H₂) as follows:NbCl₅+CH₄+½H₂

14NbC+5HCl

Coatings of ZrC, TaC or UC can be deposited by analogous reactions. Theprimary process variables that must be optimized are temperature,pressure, reactant concentration and flow rate, and deposition time.Using CVI, multiple materials may be deposited simultaneously in awell-mixed state as a single deposit. Optionally, after CVIinfiltration, exposure to high temperature hydrogen may be used toremove the underlying RVC foam skeleton, and any free carbon in thedeposited coating. Removal of the underlying skeleton using hydrogen oroxygen has virtually no impact on the mechanical performance of thefoam, since the properties are primarily determined by the stiff metalcarbide coating.

In the chemical vapor infiltration (CVI) process for a single metalcarbide, the appropriate metal in pellet form is first chlorinated andthen flowed over a heated substrate. Hydrogen and a carbon source areadded to the system. Through a combination of thermal decomposition andchemical reaction, the carbide deposits on the heated substrate surface,while HCl gas is removed from the reaction chamber by a vacuum system.Deposition of more than one metal carbide simultaneously is morecomplicated because the metal chlorides must be well mixed and in thedesired ratio in order to form a coating of the desired composition andhomogeneity. For the case of the simultaneous deposition of UC, NbC, andZrC in the present invention, one approach is to chlorinate each metalseparately and then mix the gases together prior to reaching the heatedsubstrate. This approach requires independent control of three separatechlorine sources to uniformly mix the three chlorides.

An alternate approach is to fabricate a pellet containing all of thethree metals mechanically mixed together. In this case, fine powders(e.g., −325 mesh powders, 0.0017″ diameter) are mechanically mixed in anappropriate weight ratio, e.g. 10% U:45% Zr:45% Nb and then mechanicallypressed under high pressure to create a pellet, e.g., a cylindricalpellet 0.5″ dia.×0.25″ long. The “mechanically alloyed” pressed pelletcontaining the three metals is then used in the CVI process describedabove.

A third approach is to manufacture a homogenous pellet that is a truemetallurgical alloy of the three metals. This can be done by, forexample, by liquid-phase sintering at very high temperatures. A eutecticalloy of the two or three-carbide alloys can be produced this way.

In general, fine-grained, fully dense coatings deposited by CVD havebetter stiffness and strength than do bulk materials having the samecomposition fabricated by powder processing. The elastic moduli of suchCVD films have been regularly measured up to 25% higher than those ofthe bulk materials. RVC foam is extremely well-suited as a lightweightsubstrate onto which very high-stiffness coatings may bedeposited/infiltrated by CVD/CVI. Since the modulus of the depositedfilm is so much greater than that of the vitreous carbon foam skeleton,the carbon foam has essentially no influence on the properties of thefinal product; it merely acts as a “locator” for the deposited films.Ceramic foams fabricated via CVI exhibit significantly greater thermaland mechanical fracture toughness than do monolithic ceramics since theligamental structure severely inhibits crack propagation.

Optionally, a protective coating, e.g. ZrC, may be vapor deposited ontop of the layer(s) of nuclear fuel as an additional moderator orencapsulation barrier. The protective coating can contributesignificantly to the overall porous body's strength

Other coating techniques may be used to deposit the nuclear fuel andrefractory metal carbides or carbonitrides, including chemical reactiondeposition (CRD), physical vapor deposition (PVD), electrolyticdeposition (ED), cathophoresis deposition (CD), electrophoresisdeposition (ED), and sol-gel coating (SGC), and a liquid “painting”technique that uses vacuum infiltration to draw a suspension of finepowder in a liquid binder into the porous body, followed by baking todrive off the liquid binder. Also, a “melt-infiltration” process may beused as a method of introducing the desired metals into the foamstructure and coating the ligaments, followed by conversion to atricarbide form. Also, the fuel material may be “cast” into athermally/structurally stable foam material.

FIG. 3 illustrates a schematic view of an example of a porous body withinternal structure coated by nuclear fuel, according to the presentinvention. Porous body 21 comprises interconnected ligaments 20, andinterconnected pores 16. Ligaments 20 are coated with nuclear fuel 18,which may be a continuous or discontinuous coating. Optionally, thethickness, T, of the coating of nuclear fuel 18 may be less than orequal to about 10% to 20% of the diameter, D, of pore 16. For example,if the pore diameter=800 microns, then the thickness of the coating ofnuclear fuel may be less than or equal to 80-160 microns. Alternatively,the thickness of the coating of nuclear fuel 18 may be less than orequal to 50 microns.

The internal structure (e.g., ligaments/fibers 20) of porous body 21 maycomprise one or more materials selected from the group consisting ofcarbon, graphite, Zr, Nb, Mo, Hf, Ta, W, Re, TiC, TaC, ZrC, SiC, HfC,BeC₂, B₄C, NbC, GdC, HfB₂, ZrB₂, Si₃N₄, TiO₂, BeO, SiO₂, ZrO₂, HfO₂,Y₂O₃, Al₂O₃, Sc₂O₃, and Ta₂O₅. Alternatively, the internal structure(e.g., ligaments/fibers 20) may comprise an open-celled foamstructure/skeleton comprising a carbon-bearing material selected fromthe group consisting of carbon bonded carbon fiber (CBCF) foam,reticulated vitreous carbon (RVC) foam, pitch derived carbon foam(PDGF), and graphite foam.

FIG. 4 illustrates a schematic cross-section view, Section A-A, of theexample of FIG. 3 of a ligament 20 of a porous body coated with nuclearfuel 18 and an outer barrier coating 24, according to the presentinvention. Ligament 20 may have a circular cross-section, and is coatedwith a single layer of nuclear fuel 18. Fuel layer 18 is overcoated withbarrier coating 24. Barrier coating 24 protects the underlying nuclearfuel layer 22 from exposure to hot hydrogen gases, and can serve as abarrier to prevent migration of fission products, especially fissionproduct gases, from leaving fuel layer 18 and migrating into the gascoolant stream. Barrier coating 24 may comprise one or more materialsselected from the group consisting of NbC, ZrC, BeO, BeC₂, ZrC₂, SiC,pyrolytic carbon, diamond, and diamond-like carbon. Barrier coating 24may have a thickness of about 25 microns.

FIG. 5 illustrates a schematic cross-section view, Section A-A, of anexample of a hollow ligament 20 (with hollow interior 58) of a porousbody, coated by a layer of nuclear fuel 18 and protective coating 24,according to the present invention.

FIG. 6 is a SEM micrograph magnified 15× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton before being infiltrated withnuclear fuel. The RVC skeleton has been compressed in one direction tocreate elongated pores aligned along an axis.

FIG. 7A is a SEM micrograph magnified 15× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after infiltration with TaC—NbC—ZrC.TaC was used in this sample as a substitute for UC.

FIG. 7B is a SEM micrograph magnified 45× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after infiltration with TaC—NbC—ZrC.TaC was used in this sample as a substitute for UC. The test samplesshown in FIGS. 7A and 7B were made using a mechanically alloyed, pressedpellet having the following metal ratios: 41% Zr:41% Nb:18% Ta. Then, aRVC foam skeleton was infiltrated with the tri-carbide coating (Zr, Nb,Ta)C using CVI with these pressed pellets. The results are shown inFIGS. 7A and 7B. As can be seen, the uniformity of the tri-carbidecoating (Zr, Nb, Ta)C is good. Samples were cross-sectioned and polishedand examined under optical microscope at 100× and 150×. The mixedtri-carbide coating was seen to be fully dense, with no significantporosity, and exhibited a columnar microstructure. Energy-dispersiveZ-ray spectroscopy (EDS) was performed on coated graphite and foamstructures, and the phases present were determined by X-ray diffraction(XRD). The analysis showed that is was feasible to simultaneouslydeposit the three carbides, creating a homogeneous solid solution alloyby building at the molecular level, and to infiltrate the material intoa foam structure. The XRD analysis confirmed full carburization of thethree metals, with no residual elemental metal remaining. Theinfiltrated material was nominally 5 vol % tri-carbide coating and 3 vol% RVC foam skeleton core, giving a total open porosity of 92% for theentire fuel element (with no closed porosity).

However, although tri-carbide compositions in the range of the targetwere established in these examples, the composition was inconsistentin-between infiltration runs because of the mechanically pressed pelletfeed material. As discussed previously, the pellets contained separatephases of the three metal materials. During the chlorination process,the three metals react with chlorine at slightly different rates and,hence, the uniformity of mixing of the metals within individual pelletswas imperfect. This lead to varying concentration of the three metalcarbides depositing on heated substrates from pellet to pellet. However,as stated earlier, these inconsistencies can be eliminated bymetallurgically alloying the three metals into a homogenous pellet, forexample, by liquid-phase sintering at very high temperatures.

FIG. 8A is a SEM micrograph magnified 250× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after infiltration with individuallayers of TaC/NbC/ZrC. The outer surface of ZrC appears to be continuousand crack-free. TaC was used in this sample as a substitute for UC.

FIG. 8B is an optical micrograph of a polished cross-section of alayered tri-carbide foam sample, magnified 40×. The individual layers ofthe RVC foam skeleton, TaC, NbC, and ZrC can be seen. TaC was used inthis sample as a substitute for UC.

FIG. 8C is a SEM micrograph magnified 400× of a polished cross-sectionof a layered tri-carbide foam sample. The individual layers of TaC, NbC,and ZrC can be seen. However, because this is a backscatteredsecondary-electron image (BSE) the inside core of RVC foam skeleton cannot be seen (since carbon is a low-Z material, it appears black, as doesthe empty central void). The light-colored band is the NbC layer. TaCwas used in this sample as a substitute for UC.

FIG. 9A is a SEM micrograph magnified 27× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after being infiltrated with layersof TaC/NbC/ZrC to a total of 10% vol. density. TaC was used in thissample as a substitute for UC.

FIG. 9B is a SEM micrograph magnified 400× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after being infiltrated with layersof TaC/NbC/ZrC to a total of 10% vol. density. Because this is abackscattered primary-electron image, the inside core of RVC foamskeleton can be easily identified.

FIG. 10A is a SEM micrograph magnified 27× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after being infiltrated with layersof TaC/NbC/ZrC to a total of 30% vol. density.

FIG. 10B is a SEM micrograph magnified 400× of a 65-ppi reticulatedvitreous carbon (RVC) foam skeleton after being infiltrated with layersof TaC/NbC/ZrC to a total of 30% vol. density.

FIG. 11A illustrates a schematic view of an example of a porous nuclearfuel element 50, where the interconnecting ligaments 52 of the porousbody are made of nuclear fuel 52, according to the present invention.Ligaments 52 comprise one or more compounds selected from the groupconsisting of solid-solution uranium bi-carbides, solid-solution uraniumtri-carbides, and solid-solution uranium carbonitrides. Optionally, abarrier coating (not shown), e.g., ZrC, may be applied on top of thenuclear fuel ligament 52.

FIG. 11B illustrates a schematic cross-section view, Section B-B, of theexample of FIG. 11A of a hollow, triangular-shaped ligament 52 (withhollow interior 58) of a porous body, wherein the ligament is primarilymade of nuclear fuel, according to the present invention. This structurecan be fabricated using the CVI method described above, followed byremoval of the RVC foam skeleton by baking in an oxygen orhydrogen-bearing atmosphere. Removal of the internal supportingstructure made of carbon (e.g., RVC) can help to improve neutronicperformance of fast nuclear reactors by preventing the slowing down offast neutrons by scattering from low-Z, carbon-bearing underlyingstructure in the ligaments. Additionally, a barrier coating (not shown),e.g., ZrC, may be applied on top of the nuclear fuel ligament 52.

FIG. 12 illustrates a schematic isometric view of an example of anuclear fuel element 40 comprising porous nuclear fuel 44 encased withinmetal cladding 42, according to the present invention. Gas coolant 46flows through the gas-permeable, porous nuclear fuel 44, exchanging heatwith a high heat transfer efficiency from the high-porosity nuclearfuel, due to the large extended surface area of the porous fuel, and ata high temperature due to the thinness of the nuclear fuel itself.

Referring still to FIG. 12, the composition of the porous nuclear fuel44 encased within cladding 42 may vary along the length of the element40. This axial variation in composition or porosity can be achieved byvarying the foam skeleton properties and/or the amount or type of fuelinfiltrated along the length. Additionally, neutron absorbers made ofHfC and/or B₄C may be added to the nuclear fuel 44, for example, by asimilar CVI process, to provide control of the neutron profile, forexample, along the length of cylinder 42. Shapes of the metal cladding42 other than cylindrical may be used, such as a oval, square, orhexagonal shape, as is well known in the art.

FIG. 13 illustrates a schematic cross-section view of an example of anannular porous nuclear fuel element 60 in the shape of a thick-walledcylinder 62, having a closed end 64, where gas cooling 65 enters throughthe open central core and exits by flowing radially outwards through theporous fuel 62.

Experimental Flow and Heat Transfer Tests

Six, layered TaC/NbC/ZrC foam specimens were fabricated and delivered toSandia for flow test experiments. All were 65-ppi foam cylinders ofnominal dimensions 1.5″ ID by 2.0″ OD by 2.5″ long, with tri-carbidedensities of 15%, 17%, 22%, 26%, 3%4, and 42% by volume, respectfully.In each case, the composition of the layered carbide coating was veryclose to a target of 18 vol % TaC/41 vol % ZrC/41 vol % NbC (usingtantalum as a surrogate for uranium).

Three of the six prototype foam modules supplied by Ultramet Inc. weretested at Sandia's Plasma Materials Test Facility. Flow testing wasperformed on the 85%, 78% and 54% porosity modules. Helium mass flowsfrom 0 to 35 g/s were used at an absolute background pressure of 4 MPato characterize the pressure drop across the 6 mm thick foam usingradial flow. The test results demonstrate that both the 85% and 78%porosity foams have an extremely low pressure drop, less than 350 Pa(<0.05 psia), at the highest mass flows (28 g/s) obtainable on thehelium flow loop. As expected, the highest density foam had the largestpressure drop of 24 kPa (<3.5 psia) at 27 g/s of mass flow. The 85% and78% porosity foams could be used in either radial flow for a NEPapplication or axial flow for an NTP application; or a combination ofthe two geometries is possible for a hybrid, bi-modal application.

Thermal test data were obtained on the two modules of most interest, 85%and 78% porosity. The modules were ohmically heated to between 200 and300° C. with no helium flow. After a steady state temperature wasachieved across the module, a constant helium flow rate of 0.9 g/s wasquickly started and the foam thermocouple responses were tracked duringthe cooldown. Since the mass of the foam was known, as well as thespecific heat of the tricarbide (0.329 J/gC) it was easy to determinethe power loss during cooldown. Approximately 524 W was removed from the85% module and 585 W from the 78% module during the first few seconds ofthe cooldown. Here, the conductive losses to the support structure aresmall, and radiation losses are negligible. The slope of each cooldowncurve provides a useful estimate of the effective convective heattransfer coefficient, h_(eff). Values of 489 W/m²K and 627 W/m²K wereobtained for the 85% and 78% porous modules. The area to volume ratiofor 65 ppi foam is approximately 5100 m²/m³ for the 85% porosity foamand 7100 m²/m³ for the 78% porosity foam. The average effective h_(eff)is, hence, a factor of 28 and 40 above that of a smooth tube of similardimensions or an equivalent channel in a pin-type core geometry.

The thermal tests revealed that the 78% porosity foam performed betterthan the 85% material. This is mostly attributable to the highersurface-to-volume ratio. Yet, the penalty in additional pressure dropwas minuscule. The higher density foam also has an advantage in allowingfor the use of more fissile material at lower enrichments than the 85%porous foam. Computational fluid dynamics (CFD) modeling revealed thatan 87% porous foam would have only slightly better heat transfer than a77% case using a ligament thermal conductivity of 25 W/mK. Thecalculated variation in thermal performance using two different thermalconductivities in the ligaments was shown previously. As the ligamentconductivity, k, increases to 50 W/mK, the 85% foam begins to showbetter performance than the 77% foam. However, for the k=23 W/mKligaments in these tricarbide prototypes, the higher thermal mass andsurface area of the 77% porosity foam had a slightly greater effect inthe experiment than the models predicted. It is clear, however, that thethermal performance between the 77% and 85% porous foam is very similarand either would be a good candidate for continued development. Forspace applications, cores of compact size and high power density arepreferred to minimize mass, and this pushes the design to higher densityfoams than thermal efficiency alone might dictate.

These initial experiments have established the feasibility of using atri-carbide foam material for space nuclear reactor applications.Thermal modeling of the foam at Sandia indicated a substantialimprovement in heat transfer between the high surface area foam fuel andthe coolant relative to the high delta-T present between the fuelcenterline and coolant in pin-type fuels. The overall efficiency of theheat transfer was shown to increase dramatically as the thermalconductivity of the tri-carbide ligaments is increased from 25 W/mK to50 W/mK, a level that is clearly achievable at a 2700 C operatingtemperature. A porosity range of between 77% and 85% was identified asoptimal during initial modeling work and was supported by the results ofexperimental testing.

Calculations indicate that either a small highly enriched core can beassembled that can operate at high power for brief times as required forthermal propulsion; or, a very large core of low enriched fuel can bemade that can operate at low power for very long durations as requiredfor electric propulsion systems. The results of a reactor scooping studyshowed that even 80% porous foams with modest fissile enrichments (˜20%)produce sizes that are reasonable for space applications.

In summary, a matrix of six, layered tri-carbide foam thermal/gas flowtest specimens were fabricated, demonstrating the ability to tailor thefoam density to meet application requirements. Compression testing offoam before and after 2500 C hydrogen exposure for 60 minutes confirmedboth the high strength of the foam, as well as the retention of strengthfollowing hot hydrogen exposure.

Helium flow testing performed at Sandia demonstrated the extremely lowpressure drop across the highly porous foam structures, and porositylevels were identified (78-85% porous) that could be used in eitherradial flow for a NEP application or axial flow for an NTP application;or a combination of the two geometries is possible for a hybrid,bi-modal application. The results of thermal testing revealed that foamin the lower end of this porosity range performed better because of thehigher surface-to-volume ratio, and the higher density foam also has anadvantage in allowing for the use of more fissile material at lowerenrichments.

The particular examples discussed above are cited to illustrateparticular embodiments of the invention. Other applications andembodiments of the apparatus and method of the present invention willbecome evident to those skilled in the art. It is to be understood thatthe invention is not limited in its application to the details ofconstruction, materials used, and the arrangements of components setforth in the following description or illustrated in the drawings.

The scope of the invention is defined by the claims appended hereto.

What is claimed is:
 1. A process for fabricating a porous nuclear fuelelement, comprising: a) providing a skeletal structure comprisinginterconnected pores; and b) performing chemical vapor infiltration(CVI) of the skeletal structure with a nuclear fuel material; whereinthe process of chemical vapor infiltration comprises: i) heating theskeletal structure; ii) forcing a precursor gas through the heatedskeletal structure; iii) decomposing the precursor gas inside of theheated skeletal structure; and iv) condensing the decomposed gases fromthe vapor state as a solid deposited on internal and external surfacesof the skeletal structure to form a skeletal structure comprising solidfuel ligaments; wherein the skeletal structure comprises a reticulatedvitreous or glassy carbon foam.
 2. The process of claim 1, furthercomprising: machining the skeletal structure to near-final dimensionsprior to performing chemical vapor infiltration.
 3. The process of claim1, further comprising: coating the solid fuel ligaments with a barriercoating comprising one or more materials selected from the groupconsisting of: SiC, NbC, ZrC, BeO, BeC2, ZrC2, SiC, pyrolytic carbon,diamond, and diamond-like carbon.
 4. The process of claim 1, wherein thestep of providing a skeletal structure comprising a reticulated vitreousor glassy carbon foam comprises: a) providing polyurethane foam; b)impregnating the polyurethane foam with a carbon-bearing resin; and thenc) pyrolizing the resin-impregnated polyurethane foam to produce areticulated vitreous carbon foam skeletal structure comprising vitreousor glassy carbon.
 5. The process of claim 1, wherein precursor gascomprises a mixture of niobium pentachloride, zirconium pentachloride,uranium pentachloride, methane and hydrogen; wherein the precursor gasdecomposes during chemical vapor infiltration to deposit solid uraniumtricarbide (U,Zr,Nb)C on the internal and external surfaces of theskeletal structure.
 6. The process of claim 1, wherein the stoichiometryof the solid uranium tricarbide deposit comprises (U0.1Zr0.77Nb0.13)C0.95.
 7. The process of claim 1, wherein the solid fuel ligamentscomprises a uranium tricarbide solid-solution alloy having a density ofthe alloy greater than or equal to 99% of theoretical density.
 8. Theprocess of claim 1, wherein the skeletal structure is formed ofligaments comprising one or more materials selected from the groupconsisting of: carbon, graphite, Zr, Nb, Mo, Hf, Ta, W, Re, TiC, TaC,ZrC, SiC, HfC, BeC2, B4C, NbC, GdC, HfB2, ZrB2, Si3N4, TiO2, BeO, SiO2,ZrO2, HfO2, Y2O3, Al2O3, Sc2O3, and Ta2O5.
 9. The process of claim 1,wherein the solid fuel ligaments comprise (U,Zr)C or (U,Nb)C.
 10. Theprocess of claim 1, wherein the solid fuel ligaments comprise a uraniumtricarbide alloy selected from the group consisting of (U,Zr,Nb)C;(U,Zr,Ta)C; (U,Zr,Hf)C; and (U,Zr,W)C.
 11. The process of claim 1,wherein the solid fuel ligaments comprise (U,Zr)CN or (U,Ta)CN.
 12. Theprocess of claim 1, wherein the solid fuel ligaments comprise(UWZrXNbY)CZ; where 0.04<W<0.12, 0.45<X<0.9, 0<Y<0.45, and 0.92<Z<1.0.).